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By B. J. Bayles, J. A. Ford, M. J. Salkind

Unidirecltonally sulidijied AL-A13Ni was found to exliibit excellent microstructural stability to 508°C. Above this temperature the transverse micvostructure coarsened in a manner analogous to "Ostwald Ripening", but the A13Ni whiskers did not spheroidize or foreshorlen. The coarsened whiskers reinforced the matrix as effectirely as the finer whiskevs and no reduction in room-temperature tensile strength mas noled after exposure for 90 hr at 96 pel 0-1 the absollite melling temperature of the eulectic. The slability was auribucted to the presence of low-index (and presurrlably lor-energy) cvystallographic interjaces between phases. The driving force for microstructural coursening was presumed to be the reduction of interfacial energy between the phases. Finer micvoslructures were found to coarsen more rapidly than those which weve imitally coarser. FIBER-reinforced metals have been shown to exhibit good short-time tensile properties at elevated temperatures.1-4 However, many elevated-temperature applications require long-time property retention which dictates the necessity for microstructural stability. Two types of instability are common to fiber composities. The first, chemical instability, is due to chemical reaction between the matrix and the relnforcing phase. Petrasek and weeton5 have shown that chemical reaction between tungsten wire and several copper alloy matrices resulted in degradation of the reinforcing phase and reduction in strength of the composite materials. Kreider and Leverant4 have noted similar behavior for boron-reinforced aluminum above 400°C. The second type of instability arises in systems in which the phases are essentially chemically stable with respect to each other and is characterized by spheroidization and/or agglomeration of the reinforcing phase. parratt6 has reported this type of degradation to occur relatively rapidly at moderate temperatures for nickel and cobalt alloys containing whiskers of silicon nitride, aluminum oxide, and silicon carbide and has identified it as "physicochemical" instability. Although both types of instability are the result of the differences in the chemical potential between the fibrous phase and the matrix, they may be differentiated by the fact that the second type, "physicochemi-call' instability, results in little or no net change in the composition or amount of each phase present. It does, however, result in a morphological change which profoundly affects the strengthening mechanism. Chemical interaction. on the other hand, results in a change in the composition and amount of one or more of the phases present, or the appearance of a new phase. In the latter case, the appearance of the mi-crostructure, that of a fibrous phase aligned in a matrix, may remain essentially unchanged, but the chemical degradation of the fiber results in markedly reduced composite strength. This investigation was initiated in an attempt to evaluate the microstructural and related mechanical property changes of a whisker-reinforced eutectic composite after elevated-temperature exposure. The composite system chosen for this study was the uni-directionally solidified A1-A13Ni eutectic which Lemkey. Hertzberg, and Ford7,8 demonstrated to behave as a whisker-reinforced composite. The microstructure of the unidirectionally solidified material consists of aligned needles of A13Ni in a matrix of aluminum. EXPERIMENTAL PROCEDURE Eutectic ingots having a nominal composition of 6.2 wt pct Ni were produced by melting 99.99+ pct A1 and 99.99 pct Ni in recrystallized alumina crucibles in an argon atmosphere. The analyses of the starting materials are listed in Table I. These master heats were cut up? melted by induction heating, and unidirectionally solidified vertically in 1/2 -in.-ID. 6-in.-long graphite crucibles, in a dynamic argon atmosphere. Therma1 gradients in the liquid were determined by measuring the temperature in the melt as a function of apparatus travel distance, as discussed by Kraft and Albright.9 The rate of solidification was assumed to be equal to the uniform rate at which the crucible assembly was withdrawn from the furnace. Ingots representative of nominal growth rates of 2, 5, and 11 cm per hr were produced and the solidification parameters are summarized in Table 11. Ingots used for met allographic analysis were pre-

Jan 1, 1968

By G. R. Belton, R. J. Fruehan

The Knudsen cell-mass spectrometer combimtion has been used to study the Fe-Al and Ag-Al liquid alloys. By application of the recently developed integration technique to the measured ion-current ratios, activities have been derived for the Fe-A1 system at 1600° C and for the Ag-Al system at 1340"C. The results are partially represented by the following equations: Internal consistency between the data on silver-rich and iron-rich alloys is demonstrated by application of the literature measurements on the distribution of aluminum between the nearly immiscible liquids iron and silver. The usual restrictions on the ratio of the mean free path of the escaping atoms to the orifice diameter of the Knudsen cell are shown not to be limiting in this technique. DESPITE the importance of a knowledge of the activity of aluminum in understanding deoxidation equilibria in molten steel, no direct studies have been made of activities in liquid Fe-A1 alloys at steel-making temperatures. Lower-temperature direct studies have, however, been carried out on aluminum-rich liquid alloys by Gross, Levi, Dewing, and Eilson' at 1300°C and by Coskun and Elliott' at 1315°C. Apart from phase diagram calculations by Pehlke, other determinations have been indirect and were made by measurement of the distribution of aluminum between iron and silver475 and combination of these data with extrapolated activities in the Ag-A1 system.~-% ecently, however, Woolley and Elliott have made a significant contribution by directly measuring heats of solution in the Fe-A1 system at 1600°C. The present authorslo have recently employed a Knudsen cell-mass spectrometer technique in a study of activities in iron-based liquid alloys. In this technique activities and heats of solution are determined from a series of measurements of the ratio of ion currents of the components; and since ion-current ratios are used, problems caused by changes in instrument sensitivity or cell geometry are overcome. Results obtained for the Fe-Ni system were found to be in excellent agreement with previous work, thus demonstrating the reliability of the method. The present paper describes a similar study of activities in the liquid Fe-A1 and Ag-A1 systems, this latter system being included in order that a meaningful comparison can be made with the above-mentioned indirect studies. INTEGRATION EQUATIONS A detailed derivation of the equations used to determine the thermodynamic properties from the measured ion current ratios has been given elsewhere;'' however it is useful to summarize them here. By the combination of the Gibbs-Duhem equation with the direct proportionality between ion-current ratios and partial pressure ratios, it was shown that for a binary system at constant temperature and pressure: where al is the activity of component 1 with pure substance as the standard state, N, is the atom fraction of component 2 in the solution, and I; and t'2 are ion currents of given isotopes of the components. The activity coefficient is given by: this latter equation being more suitable for graphical integration. Combination of Eq. [l] with the Gibbs-Helmholtz equation gives an expression for the partial molar heat of mixing: EXPERIMENTAL A Bendix Time-of-Flight mass spectrometer model 12! fitted with a 107 ion source and a M-105-G-6 electron multiplier, was used to analyze the vapor effusing from the Knudsen cell. The arrangement of the Knudsen cell assembly was essentially that of the commercial instrument (Bendix model 1030) but with several modifications. Instead of heating with a single tungsten filament, a cylindrical tantalum-mesh heater was employed. Up to 1400°C simple resistance heating was used but above this temperature electron bombardment between the tantalum mesh and the tantalum cell susceptor was necessary. The temperature was measured by means of a Leeds and Northrup disappear ing-filament type optical pyrometer sighted on an essentially black-body hole in the side of the cell. Details of the temperature control, temperature measurement, and in situ calibration of the optical pyrometer can be found elsewhere.I0 In the investigation of the Fe-A1 system the Knudsen cells were constructed of thoria crucibles with fitted thoria lids (Zircoa). The cells employed in investigating the Ag-A1 alloys were made up of high-purity alumina crucibles (Morganite) with lids of recrystal-lized alumina (Lucalox). The cells were 0.370 in.

Jan 1, 1970

By K. G. Davis

Dilute alloys of silver and of thallium in tin have been solidijzed unidirectionally under controlled conditions, to study the segregation associated with a cellular interface under conditions where both thermal and solute convection are present. Autoradiography and radioactive tracer counting techniques were combined with electron-probe microanalysis to study both macro- and microsegregation. It was found that, for concentrations giving only small amounts of constitutional supercooling, cell formation had little effect on the macroscopic distribution of solute along the specimen. At higher concentrations the effective distribution coefficient was higher than that expected for a smooth interface. Node spacing was independent of initial solute content at lower concentrations, becoming greater as keff increased. Silver content at the segregation nodes of silver in tin alloys was independent of initial concentration and considerably in excess of the eutectic composition. SINCE the investigation of cell formation at advancing solid-liquid interfaces by Rutter and Chalmers,' a large volume of work has been dedicated to the determination of solidification conditions under which a planar interface will break down into cellular form. Early experiments were explained satisfactorily by the concept of constitutional supercooling,2 but, due to poor measurement of temperature gradients in the liquid, lack of accurate data on liquid diffusion and equilibrium distribution coefficients, and uncertainty about the effects of thermal and solute convection, these experiments cannot be used as proof for the theory. More recent work, however, has shown that under conditions where convection is eliminated or can be ignored good correlation is observed.3,4 Investigations into segregation at cell caps5 and at cell nodes6-'' have been made, but no measurements appear to have been done on the overall, macroscopic segregation down a unidirectionally solidified rod of material which has solidified with a cellular substructure. This has practical importance in casting, where regions of material with cellular substructure are often encountered, and also in zone refining where the thermal conditions necessary for a planar interface are unattainable. Further, as will be shown, the macroscopic segregation can give information on the following question. Granted that a cellular solid-liquid interface develops from a planar one when the conditions for constitutional supercooling are exceeded, how much supercooling is present after the cells have formed? EXPERIMENTAL PROCEDURE AND RESULTS Specimen Preparation. Specimens 25 cm long with a square cross section 0.6 by 0.6 cm were grown in graphite boats by solidification from one end. Alloy compositions are given in Table I. Two specimens of each composition were grown. The tin was 5-9 grade and the silver and thallium both 4-9 grade. Ag110 and Tl204 were used as tracers. Each composition had the same quantity of tracer so that auto radiographs of specimens containing different concentrations of the same element could be easily compared. Thermocouples inserted through the lid of the boat into a dummy specimen showed that, over the first 10 cm of growth, thermal conditions were quite steady, with a rate of interface advance of 5.8 cm per hr and a temperature gradient in the melt ahead of the interface of 3.0°C per cm. The specimens were seeded from tin crystals of a common orientation to eliminate orientation effects. Dilution of the specimen by seed material was minimized by the provision of a narrow neck between specimen and seed crystal. Macrosegregation. After growth, the specimens were sectioned with a spark cutter. The rods of silver alloy were cut into 1-cm lengths and analyzed for Ag110 using a y -ray counter with fixed geometry. The specimens containing thallium were cut into 2-cm lengths and analyzed for T1 204 by taking 13 counts from each end of the cut lengths through an aperture in lead sheet approximately 0.4 cm square. The results are summarized in Figs. 1 and 2. To find the effective distribution coefficient for the silver in tin alloys under smooth interface conditions, the region of substructure at the bottom surface of one of the 10 ppm specimens, see Fig. 3, was removed by spark machining before counting. Autoradiography. For both alloy systems the samples were polished on sections taken alternately parallel and perpendicular to the growth direction, and autoradiographed by placing the polished surfaces in contact with Kodak "Process Ortho" film. Figs. 3 and 4 show the structures revealed. The alloy containing 10 ppm Ag showed substructure only after a few centimeters of growth, and then substructure was limited to a narrow layer at the base. The "speckled" substructure reported previously in this system4 is here clearly seen to be an intermediate stage between planar and cellular interface conditions. The other samples show a remarkable similarity considering

Jan 1, 1969

The organization of the American Institute of Mining Engineers, 75 years ago, parallels the beginning of present-day steel-producing methods in the United States. This early association with the industry is emphasized by the fact that the first open-hearth furnace constructed and operated in this country was then only three years old, and the first commercial bessemer converter was just seven years old. Total steel production in the year 1871, as recorded by the American Iron and Steel Institute, was 84,000 net tons, of which 2,000 tons was produced by the open-hearth process, 45,000 tons by the bessemer, and 37,000 tons by the crucible, puddled iron, and other processes. The bessemer process had thus become the major source of steel production. Its capacity expanded progressively during the following decade, and by 1880 was the source of approximately 86 pct of total steel production. Most of this tonnage was made in eastern Pennsylvania and used in rail manufacture. At about this time, however, important changes began to occur which gave impetus to the production of steel on a much larger scale. The development of Lake Superior ores from 1870 to 1890 and the exploitation of the Mesabi Range in 1892 provided the industry with great quantities of high-grade ores. The abundance of coking coals and limestone in the region adjoining western Pennsylvania and the development of cheap water transportation on the Great Lakes tended to shift the center of iron and steel manufacture to an area west of the Allegheny Mountains. Encouraging these new develop-

Jan 1, 1948

Application of technology to the search for new deposits went on apace in 1963. Traditional methods, aided by modern communications, were successful in some out-of-the way corners of the world that had not been adequately prospected, but the co-ordination of several mutually supporting techniques is necessary to define exploration targets in most areas today. Although these obviously include geophysics and geochemistry, which are treated elsewhere in this issue, geology continues to play a major role. As the search focuses on less and less obvious targets there will be more need than ever to "pull out all the stops." Any anomaly may have significance and must be evaluated as a possible lead. Ability to detect them increases as the laboratory workers refine techniques for measuring physical, chemical, and isotopic properties; it will be increasingly the responsibility of the exploration geologist to understand and apply these to economic problems. Simple propecting, the search for outcrops of recognizable ores of traditional types, was perhaps most successful for iron. Potentially large deposits of high-grade magnetite-specularite-hematite ore have been found in the Mary River area near the north end of Baffin Island. The ore holds up four hills that stand some 1500 ft above the surrounding country; obviously they were ripe for discovery by the first person to come across them, in this case Murray Watts, president of British Ungava Explorations, Ltd, As preliminary examination indicates a lump ore averaging 68 to 70% iron, these deposits may be commercial in spite of the unfavorable climate and their isolation from world markets.

Jan 2, 1964

C. M. Bishop (Aircraft Div., A. V. Roe Cnnada Ltd., Toronto)—Reference is made to the coherency hardening of ß due to ß' and also to the hardening of ß due to a coherent precipitation of the a phase. What is meant by the term "ß'?" Would not the coherency hardening of ß during cooling be due to precipitation of a sub-microscopic dispersion of a, as is the case during reheating a ß-quenched alloy? I feel that these two phenomena are essentially the same mechanism, i.e., precipitation of a from unstable or supersaturated ß. From the text, it would appear that Tables III and IV are intended to show the effect of the martensite transformation on the properties of Ti-Mn alloys. Referring to Table 111, the ß-quenched alloys would contain a greater proportion of retained ß than would the equilibrated alloy. Could not the increase in properties be attributed to this, along with possible coherency hardening of the ß? In the same way, the structures referred to in Table IV contain ß from which they could derive their hardness. It would appear to be difficult, if not impossible, to assess the role of the martensite transformation from the data contained in these tables. F. C. Holden, H. R. Ogden, and R. I. Jaffee (authors' reply)—In reply to the first question raised by Mr. Bishop, the term "ß"' designates an apparent retained-p structure which is harder than it is after the most rapid cooling from the ß field. It is thought to be the result of a coherent precipitation of a. The mechanism by which hardening takes place is undoubtedly the same as that involved in the early stages of aging of a re-tained-p structure, and we are in agreement with Mr. Bishop on this point. The term "ß'" to designate the anomalously hard retained-ß structure was agreed upon at a symposium in 1953." The authors do not agree that the quantity of retained ß and/or ß' in the ß-quenched specimens of Table III is necessarily greater than that of the 600°C equilibrated specimens. The ß-quenched specimen has a structure which appears to be completely a' formed by the martensite reaction. It is recognized that there could be some retained ß (or ß') between the a plates if the M, temperature is below room temperature. This is unlikely for the 3.4 pct Mn alloy where the M, is known to be at about 640°C. Therefore, the increase in strength cannot be attributed to the presence of a greater quantity of coherency-hardened ß, but, rather, the increase in strength must be attributed to the transformation products formed, e.g., a' (martensiti-cally formed a). Concerning the data in Table IV, we agree that it is difficult to estimate the role of the martensite reaction, since the 6.4 pct Mn alloy was retained ß in the quenched condition.

Jan 1, 1955

By M. C. Udy, F. W. Boulger

AN incomplete phase diagram for the U-Ti systern was determined earlier 1 and more recently, a tentative diagram was presented for the uranium-rich end of the system.' In the present re-examination of the whole system of U-Ti alloys, high purity materials were used. Melting stock for the alloys was high purity uranium, containing about 0.09 pct C as the only appreciable impurity, and high purity iodide-process titanium purchased from New Jersey Zinc Co. Both metals were cold rolled to about 1/6 in. thickness, sheared to about I/' in. squares, and cleaned by pickling. The alloys were arc melted under a helium atmosphere in a water-cooled copper crucible. A thoriated-tungsten electrode was used. The furnace chamber was evacuated, then flushed with helium, prior to each melting. It was finally filled with stagnant helium at one atmosphere pressure. Each alloy was remelted three times after the original melting, to insure homogeneity. The alloy button was turned bottom side up before each re-melting operation. Some 22 alloys were examined. Their compositions were spaced at appropriate intervals between 100 pct Ti and 100 pct U. Analyses were made on chips taken after fabrication. The major contaminant was carbon, which varied from 0.03 to 0.08 pct. It appeared in the microstructure as titanium carbide. Alloy compositions were calculated to a carbon-free basis for consideration on the diagram. Tungsten and copper, possible contaminants from the melting operation, were generally less than 100 parts per million each. Fabrication All alloys were forged and rolled to bars approximately V8 in. square. They were clad either in SAE 1020 steel or in a 5 pct Cr-3 pct Al-Ti-base alloy, depending on the fabrication temperature. A temperature of 1800°F (980°C) was used for alloys near the compound composition. This necessitated using the titanium-base alloy, since iron reacts with titanium at this temperature, producing a low melting alloy. Other alloys were fabricated at 1450°F (790°C), using steel jackets. No iron-titanium reaction occurred at this temperature. The jackets were welded in place in an argon atmosphere. Those alloys sheathed in steel were declad and then reclad between rolling and forging operations. On the other hand, those clad with the titanium alloy were cut to a roughly rectangular shape prior to clading and were then carried through both the forging and rolling operations without opening. Those alloys near the compound composition were found to be cracked when the clading was removed. The cracked materials had been plastically deformed, however, and at least some of the cracking had OCcurred during cooling. Heat Treatment The rolled bars, after being declad and shaped to remove surface contamination, were all given an homogenizing treatment of 160 hr at 2000°F. (Samples were taken for analysis following the declading and shaping operations.) All were heat treated at the same time in one furnace, but each was sealed in a purified argon atmosphere in an individual Vycor glass tube. Argon pressure was such that it was approximately atmospheric at temperature. One end of each tube contained titanium chips and this end was heated to 1200°F (650°C) for 10 min prior to the heat treatment. This purged the atmosphere of residual reactive gases. The balance of the tube was warmed during the purge to liberate adsorbed moisture and gases, which also reacted with the hot chips. The bars were furnace cooled from the homogenization treatment. Specimens of each alloy were water quenched after 2 hr heating at 1000°, 1200°, 1400°, 1600°, 1800°, and 2000°F (540°, 650°, 760°, 870°, 980°, and 1095°C). In addition, some were treated at intermediate temperatures of 1300°, 1500°, and 1700°F (705", 815", and 925°C) and at 2150°F (1175°C). Specimens, about '/s in. cubes, were cut from the bars, sealed in individual Vycor tubes, and heat treated as described. All specimens heat treated at the same temperature were processed together. Samples were quenched by breaking the Vycor tube rapidly under water. Metallographic Examination Specimens were mounted in bakelite and ground wet on 180 grit paper held on a 1750 rpm disk. They were then ground wet by hand, using 240, 400, and 600 grit papers. The rough grinding was continued long enough to get well below the surface. Specimens were mounted separately because of the variation in the rate of etching between alloys. The specimens were polished with rouge on a 4 in., 1725 rpm wheel covered with Miracloth. Alloys on the titanium side of the compound composition were etched with a solution of 2 pct hydrofluoric acid in water saturated with oxalic acid. A few crystals of ferric nitrate were added as a bright -ener. Specimens were immersed 5 sec, polished to remove the etch, then re-etched. With the higher titanium alloys, it was often necessary to start the etch on the polishing wheel, because of the formation of a passive film. In some instances, a plain 2 pct hydrofluoric etch was satisfactory. For the alloys on the uranium side of the compound, a distinction between the compound and the uranium phase developed after standing a short time in air. This could be hastened by the application of heat, such as obtained by placing the specimen on a radiator. A deep etch was necessary to develop details in the uranium-rich phase, such as the Widmanstaetten pattern sometimes obtained by quenching y uranium. A 2 pct hydrofluoric acid solution was used for this deep etching.

Jan 1, 1955

By George S. Rice

Those who for the first time see, at a mine, a great hole caused by subsidence; or, going underground, see an extensive fall of roof or hanging wall are apt to regard such an occurrence as an accident and not a normal condition of exhaustive mining. Moreover, the public is inclined to believe that a cave or subsidence of the surface is the result of carelcssness in mining or the rapacity of the mine operators in "robbing pillars," a tern too loosely used in descriptions of coal mining. This belief is fostered by the common-law requirement that an owner of the surface or of surface improvements is entitled to full support of the surface. Court decisions have upheld this law even when that owner has sold or leased the underlying mineral, unless a clause in the sale or lease contract exempts the mining company from paying the owner of the surface for damages that may result from mining. Another reason for this belief is the earlier practice (which is still extensively followed in this country) of leaving all coal pillars unmined, except as they have been thinned by so-called robbing; by not extracting the pillars, the surface may be supported by them for years, perhaps for generations. Inasmuch as the practice of leaving permanently in the mines from one-third to one-half of the mineral is wrong, with respect to conservation of a national resource, it behooves the mining engineers to assemble data and disseminate information on the subject, pointing out that subsidence must follow complete extraction, unless expensive methods of back-filling are employed, such as are used in Germany and France. These methods inevitably increase the cost of mining and the price of the mineral produced. It is popular to speak of "waste in mining" without a realization that much so-called wasting is unavoidable unless subsidence is accepted and provision made therefor where the surface is valuable, or else the only present known alternative of hydraulically placing sand or crushed stone filling in the excavations, at an increased cost of the mineral product to the consumer.

Jan 1, 1923

By P. Chiotti, G. R. Kilp

Thermal, metallographic, vapor pressure, and X-ray data were obtained to establish the phase diagram for the zinc-zzrconiz~m system. Five compounds corresponding to the stoi-chiometric formulas ZrZn, ZrZn,, ZrZn,, ZrZn,, and ZrZn14 were observed. All these compounds, with the exception of ZrZn2, which melts congruently at 1180°C under constrained zinc-vapor conditions, undergo pexitectic reactians. The temperature at which the zinc vapor pressure is I atm for a series of alloys was determined from vapor-pressure measurements. The data obtained are summarized in the construction of a I-atm-pressure phase diagram and a phase diagram corresponding to a pressure of less than 10 atm. THE purpose of this investigation was to establish the phase diagram for the zinc-zirconium system. Thermal, metallographic, vapor pressure, and X-ray data were employed in determining the phase regions. Partial investigations of this system have been conducted by Gebhardt1 and Carlson and Borders.' Carlson and Borders studied the high-zirconium region and established the existence of a eutectic at 69 wt pct Zr with a melting point of 1015°C. The terminal phases of the eutectic horizontal were shown to be an intermetallic compound ZrZn and a solid solution of ß zirconium containing 21 wt pct Zn. The ß solid solution decomposes into ZrZn and a zirconium at 750°C. The eutectoid composition is given as 15 wt pct Zn, and the solubility of zinc in a zirconium at temperatures below 750°C is indicated to be negligible. Gebhardt studied the zinc-rich region and observed a lowering of the melting point of zinc from 419.5" to 416°C and temperature horizontals at 545" and970°C. Some preliminary observations by Chiotti, Ratliff, and Kilp were reported by Hayes.2 pietrokowsky3 has reported the compound ZrZn2 to have a cubic MgCu2 structure with ao = 7.396A. MATERIALS AND EXPERIMENTAL PROCEDURES The metals employed in the preparation of alloys were Bunker Hill slab zinc or Baker analyzed reagent granulated zinc, both 99.99 pct pure and hafnium-free iodide-process crystal bar zirconium obtained from the Westinghouse Electric Corp. The zirconium contained 200 ppm Fe, 200 ppm Si, 100 ppm C, and minor amounts of other impurities. The zirconium was milled or machined into thin chips or shavings. These were cleaned with a nitric-hydrofluoric acid solution, rinsed with water, and acetone, and dried just prior to their use in alloy preparation. The granulated zinc was similarly cleaned using dilute nitric or hydrochloric acid. Weighed quantities of these materials, 20 to 30 g total, were mixed and pressed at 20,000 to 70,000 psi to give relatively dense compacts. During the early part of this investigation the pressed compacts were placed in MgO-15 wt pct MgF, crucibles which were then sealed inside of quartz ampules. The compacts were given various prolonged heat treatments prior to their use for thermal analyses, or vapor-pressure measurements. Because of expansion of the compacts and the relatively high zinc vapor pressure it was difficult to heat to the melting temperatures of the alloys without failure of the quartz ampules. Homogenization at temperatures below the melting temperature gave brittle, porous alloys unsuitable for metallographic examination. It was also difficult to prevent condensation and segregation of zinc on the colder parts of the quartz ampules during heating and cooling operations. These problems were eliminated to a great extent by the use of tantalum crucibles. Tantalum proved to be a satisfactory container with little or no reaction between the alloys and the tantalum. Small tantalum thermocouple wells were successfully welded in the bottom of these crucibles. Pressed compacts were sealed inside the tantalum crucibles by welding on preformed caps under an argon atmosphere. Heat treating and differential thermal analysis were combined into a single operation. The experimental sample assembly is shown in Fig. 1. This assembly was enclosed inside a stainless-steel tube heating chamber which could be evacuated and filled with an inert gas. The thermocouple leads were brought out of the heating chamber between two rubber gaskets used to provide a vacuum seal for the water-cooled head. Most of the compounds in this system undergo peritectic decomposition. After heating above the temperature of a particular peritectic horizontal the sample was cooled to just below the peritectic temperature and held at temperature for several hours. The sample was then reheated through the peritectic temperature and the size of the thermal arrest, if still present, compared with the one previously obtained. If the thermal arrest was not characteristic for the alloy composition being investigated its magnitude diminished and repeated cycling and annealing eventually eliminated it. The peritectic thermal arrests characteristic of a particular composition were established in this manner.

Jan 1, 1960

By R. G. Wayland

SINCE the installation at Kropfmuehl, Bavaria, of a modern flotation concentrator in 1938, the flake and fine graphite from the Passau area can now be delivered in about any normal specified carbon content of any size range up to a flake averaging about 0.7 mm. The graphite finds a wide German and export market for crucible manufacture, pencil leads, dry cells and other uses. A controversy over the origin of the graphite deposits is being resolved in favor of syngenesis rather than epigenesis. The syngenetic theory is newly supported by the yet unpublished work of Hartmann of the Bavarian Geological Survey. Development work and exploration for graphite in the area may be changed in direction as the syngenetic theory is accepted. Crystalline graphite is produced in the area east of Passau near the junction of Germany, Austria and Czechoslovakia, as shown in Fig. 1. This is the only graphite area of importance in Germany, and Gra-phitwerk Kropfmuehl AG is the only operating firm in the area. The plant and mine are located at Kropfmuehl near Hauzenberg, Kreis Wegscheid, about 10 miles east of Passau. A narrow gage railway from the mine connects with the German Railway at Schaibing Bahnhof. The Pfaffenreuth mines date from about 1730. Until the early 20th century mining operations were carried out in a haphazard fashion. During World War I graphite mining and milling was increased, since it had to cover almost all of the crucible needs of the Central Powers. Between the wars some 11 mines were operated by two large and several small companies, but under the Nazis these were consolidated by 1938 into the Kropfmuehl enterprise. Kropfmuehl built a modern flotation mill to treat its own ores and small amounts of custom ores and tailings from the area. Since Graphitwerk Kropfmuehl AG was an I.G. Farbenindustry subsidiary, it has been under Military Government Property Control and probably will be sold to private German capital. Geology The country rock of the graphite area is part of the "kristallines Grundgebirge," the series of old gneissic and schistose rocks that constitutes the bed rock of most of the Bohemian basin and rims the Sudeten-land. The gneissic rocks of the graphite area are considered to have been metamorphosed during the Carboniferous period. They are bordered on the north by granite stocks and penetrated by numerous smaller granite and pegmatitic intrusive rocks, as shown in Fig. 1. The gneiss is classed as a micaceous, coarse-grained cordierite gneiss by most investigators. It is much metamorphosed by the granite, particularly in the north near the larger granite bodies. Interbedded in the gneiss are the graphite seams and lenses, and also beds and lenses of crystalline limestone containing disseminated graphite in noncommercial quantities. The gneiss, together with the included graphite and limestone seams and lenses, is cut and displaced by a number of granite sills of medium to fine grain and by a large number of diorite lamprophyre dikes and a few syenite-pegmatite dikes. The lamprophyre dikes are of various mineral compositions and textures, but many are banded and richly impregnated with pyrite; while the syenite-pegmatite dikes are coarse-grained with good crystals of titanite, pyroxene, uralite and other green amphiboles. Most investigators and the miners speak only of diorite and granite dikes cutting the graphite seams. The diorite dikes are later than the granite and some of the faulting, as is evident from Figs. 1 and 2. Individual graphite seams and lenses may be mined for thicknesses of several feet up to several scores of feet, and for distances of several hundred feet. The aggregate thickness of a series of some 20 seams of graphite, limestone and interbedded gneiss at Kropfmuehl is stratigraphically about 450 ft. Laterally, the graphite in a seam may pinch out or grade into crystalline limestone. Graphite crystals also are found disseminated in the gneiss itself, although in unmineable concentrations. The faults that cut the graphite seams carry graphite for some feet or tens of feet away from the seams, apparently mechanically. Similarily, the graphite lenses themselves often contain mechanically introduced inclusions of wall rock, probably from flowage during folding. Graphite crystals make up 10 to 30 pct of the fresh, mineable graphite lenses at Kropfmuehl, averaging about 20 to 25 pct after hand-sorting by the miners. In weathered lenses, the graphite concentration is said to be as high as 50 pct. The associated primary and hydrothermal minerals are dom-inantly feldspar and calcite, plus quartz, pyrrhotite, pyrite, biotite and occasional garnet, hornblende, sphalerite and galena. Associated secondary minerals include kaolin, nontronite, mangano-oxide-silicates (mog), adularia and chlorite. The superimposed suite of siliceous cementation minerals present consist largely of opal, chloropal, chalcedony, jasper, and hyalite. The kaolin is of special interest since it too was mined as early as 1730 and was used in the well-known Nymphenburg porcelain from 1756 on. The kaolin is derived from the gneiss and the syenite pegmatites. The crystals of graphite vary in size within a given seam, but in seams more than a mile away from the granite on the north the average crystal-linity is less coarse, lowering the commercial value. The lenses in the Kropfmuehl-Pfaffenreuth area are the most developed, and are the only ones with deep workings now accessible. Other similar crystalline graphite lenses are known from older workings at Habersdorf, Oberoetzdorf, Ficht, Diendorf, and

Jan 1, 1952

By W. F. Hower, W. Brown

Large-scale sand consolidation tests were conducted in an effort to determine the reasons for the successes and failures of this method of sand control. Several different consolidating materials were used in treating both clean and bentonitic sands that were packed in a chamber having a capacity of 3.3 cu ft. The results were essentially the same for all of the different consolidating materials, The data show that low-viscosity consolidating materials pumped at a relatively slow rate gave the best results. Where the formation has produced sand, the treating fluids can compress the formation, thus permitting the channeling of fluids to another horizon. Pressure-packing these zones before attempting to consolidate is recommended. Sands containing more than 4 per cent of water-swelling clays are not good candidates for consolidation. It is indicated that loose sand, particularly when it is bentonitic, can be fractured during the placement of the treating fluids. INTRODUCTION Sand production in oil and gas wells has plagued the industry for many years, and numerous cures for this problem have been suggested. Most methods have been successful to a certain degree, but the great variety of well conditions that exist in the different areas has magni- fied the problem and limited the successful use of the various systems. Four review papers1-4 present a wealth of information concerning the degrees of success that have been obtained by the different sand-control methods. The bridging of sand grains by the use of gravel packs and screens has been quite successful. However, these methods do not leave the casing clear for all types of multiple completions, and the cure does not last for the production life of the well in some instance:;. The control of loose sands by sand consolidation with resins has never been as successful as desired. It has always been hoped that such a treatment would eliminate all sand problems for the life of the well, but. initial applications, starting in the middle 1940's, were only moderately successful. Lott, et a1,3 reported a success ratio of approximately 50 per cent and made the following conclusions. The highest percentage of successes were obtained where: a. Consolidation of a zone was made at the time of initial completion or prior to the production of sand. b. The interval treated was less than 12 ft in length. c. Between 30 and 50 gill plastic/ft of producing interval was displaced through the perforations. REASONS FOR SAND CONSOLIDATION FAILURES Our own experiences in the field of sand consolidation point toward the following conditions as the major reasons for the failure of sand consolidation attempts. 1. Mud-plugged perforations and mud invasion of the formation. 2. Sand in the casing covering all or part of the perforations. This sand could be either formation sand or one of the coarser sands used as propping agents in hydraulic fracturing. 3. Holes in the casing. 4. Channels behind the casing. 5. Attempting to treat too long a perforated section. 6. Too high a percentage of water-swelling clays in the formation. 7. Formations that have produced sand. Recent attempts were made to treat perforated sections ranging from 10 to 30 ft, in wells that have produced sand, by using a straddle packer that was raised and lowered through the perforations as the consolidating material was being pumped. In most instances, the pressure required to pump fluid into the formation varied considerably as the tool was raised and lowered. This suggested the possibility that significant differences in permeability were present or that only part of the formation had produced sand. There were times when a sudden break in pressure indicated that a fracture was being formed. Research conducted several years ago concerning the problem of the control of water in air and gas drilling indicated that shale sections could be fractured quite easily. In addition, it was determined that it was easier to pump fluids into shale bodies by fracturing the shale itself, or the interface between the shale and sand, than to pump into a fluid-saturated formation. Formations that produce sand are usually adjacent to shale bodies and frequently have shale streaks of various thicknesses inter-bedded in the sand. Therefore, where shale is exposed to fluid pressure it

By R. I. Jaffee, F. C. Holden, H. R. Ogden

The effects of ternary and quaternary additions of aluminum, oxygen, and carbon on the mechanical properties of high-purity titanium-molybdenum alloys were studied for several microstructural conditions. Heat treatments were designed to produce 1) ß-quenched, 2) equiaxed ß) transformed (acicular) and 4) stabilized -P microstructures. Tension, impact, and hardness tests were performed. Transformations from the ß phase were followed by hardness measurements and metallography. The ß-to-w transformation was slowed by aluminum additions; carbon and oxygen additions increased the transformation rate. In addition to solid-solution strengthening, these solutes may increase or decrease strength by altering the /3 transformation kinetics. Aluminum additions promote the formation of strain-induced martensite, and lower yield strengths at compositions near 12 pct molybdenum.. a-P alloys are more ductile with equiaxed than with transformed microstructures. High strength levels can be reached by aging in many cases. Strength levels of the aged specimens usually are lowered slightly by straining prior to aging. ALLOYS of molybdenum with titanium form an iso-morphous alloy system which has been the subject of several investigations. Additions of substitu-tional and interstitial -stabilizing elements provide means of improving mechanical strength and con- trolling the heat-treatment response of the binary alloys. The work reported here forms a part of a continued study of the effects of microstructure on the mechanical properties of high-purity titanium alloys. Previous papers have considered the various binary alloys (Ti-Mo, Ti-Al, Ti-C, and Ti-0) that form the bases for the ternary and quaternary alloys included here. The compositions were selected to correspond with the previous binary alloys. Further, the high purity of the alloying components was maintained, so that direct comparisons with earlier work can be made. The total research program included a large number of compositional, microstructural, and testing variables. The data reported here have been selected to provide a reasonably complete and representative summary of the results. The reader is referred to the original Air Force Report for a complete tabulation of the results. EXPERIMENTAL PROCEDURES All the alloys used in this work were prepared from high-purity electrolytic titanium produced by the U.S. Bureau of Mines. A typical furnished analysis showed the following impurities in wt pct, balance This analysis is equivalent to that of iodide-refined titanium. In addition, hardness, tensile, and impact tests made using unalloyed electrolytic titanium showed it to have properties equivalent to those obtained with iodide titanium and suitable for these studies of high-purity base alloys.

Jan 1, 1962

By H. F. Silliman, Daniel R. Hull, John R. Freeman

IN the normal operation of a brass-rolling mill, sheet and strip has, for the most part, been finished in comparatively thin gauges, involving a substantial amount of coldwork and a considerable number of anneals. This condition has been disturbed by the prep0nderance of cartridge brass, which is not rolled to thin gauges, and particularly by brass for artillery cartridge cases, on which the mill may finish its work at a thickness between 0.400 and 0.900 in. Obviously, the amount of rolling and the number of anneals to which such a product can be subjected are small compared with strip that is carried to a thickness of a few hundredths of an inch. As a result, critical examination into the internal condition of metal has been made on an extensive scale at a stage of manufacture that would not, in normal times, have been chosen for complete inspection. Under these conditions, a type of defect has been recognized that is not seen at thinner gauges. It might be more accurate to call it the appearance of a defect, since it refers to an apparently laminated, discontinuous structure, which, in fact, does not exist as such, and disappears after suitable treatment. For this reason it has been called a phantom lamination, which the authors consider an appropriate term. It is the purpose of this paper to describe the phenomenon and the conditions of its occurrence and, at least tentatively, to discuss its nature and cause. Great numbers of phantoms may be found in 2: 1 brass, either hot-rolled or cold-rolled. They are also found in 70-30 brass, sometimes when hot-rolled and generally when cold-rolled. It is not believed that phantoms are confined to these alloys but this paper is limited to them for the purpose of description. HOT-ROLLED 2:I BRASS For illustration, therefore, a typical cake of hot-rolled 2: 1 brass has been chosen. This was rolled to a slab 7/8 in. thick, from a casting 5 in. thick by 24 in. wide by 18 in. high, which was believed to have been sound. After hot-rolling and cooling, the slab was sheared and a portion of such a shear cut is shown in Fig. 1. It would be difficult to convince anyone viewing this specimen that the metal was not unsound. The impression would be confirmed by the appearance of a tensile break, as shown in Fig. 2. This is an ordinary tensile-test piece from metal just described. Although there was nothing in its tensile strength or elongation that suggested unsoundness, it appeared run through with pipes and laminations. To question the reality of such a condition would imply questioning the obvious. But, considering the history of the metal, it w0uld be difficult to believe that it could be so bad as it appeared. It seemed reasonable, though, that if the defects were what they appeared to be, they could be traced back from the sheared or broken face.

Jan 1, 1945

By T. I. Moore, L. A. Painter

THE electrolytic zinc plant of the American Zinc Co. of Illinois was described by Davidson' in 1944. Since then, improvements as well as expansion of the plant facilities have been made. In order to increase the production of high grade zinc which was needed for war purposes, an expansion program designed to double the slab zinc capacity was started in 1942 and completed in March 1943. This expansion was propagated by a contract between the American Zinc Co. of Ill. and the Defense Plant Corp. The contract included the facilities of the Fairmont City, Ill., property of the American Zinc Co., where a suspension-type roaster with contact acid plant, cadmium distillation furnace, Waelz oxide and densifying plant, and horizontal retort furnaces were installed. The expanded Monsanto, Ill., plant and the additional facilities of' the Fairmont plant were designed to integrate the metallurgical treatment of zinc concentrates for the production of special high grade zinc at Monsanto with the production of acid, cadmium, high grade zinc from furnace skimmings and the Waelz treatment of leach residue at Fairmont. In general, the original flowsheet was not changed, except for the addition of the filtering, drying and reclaiming of leach residue, and the treatment of purification cake for the recovery of copper, cadmium sponge, and zinc. Fig. 1 is a flow diagram of present operations. The original plant facilities, desi-gned for 50 tons daily production of slab zinc, had some units which were more than adequate. Therefore, in expanding the facilities to 100 tons per day, it was not necessary to double all operating components. Table I gives the comparison of the changes made in the unit operating components for the original facilities, 1941, the 1943 expansion, and the 1951 facilities. During the past 11 years a number of improvenients have been made resulting in: 1—an increase in slab production, 2—higher recoveries on the calcine treated, 3—better quality of slab zinc produced, 4—higher current efficiencies, and 5—less man hours Table I. Changes in Operating Facilities Operating Unit 1941 1943 1951 Calcine unloading (pneumatic), 10 tons per hr 12 calcine unloading track hpr. and elev., 60 tons per hr 1 Calcine storage, tons 1,000 2,000 2,000 Leach tanks, 35 vol. tons. No. 3 5 6 slurry mixing 6x6 ft stainless tank, No. 1 Ball mill. 4.5 rt x 16 in. conical. No. 1 CLassifier duplex, No. 1 1 Thickeners. 50 it diam. No. 2 9 2 Filter thickeners, sq ft '-- Moore filters, sq ft 5.760 11,520 Drum filters, 10 ft diam x 16 ft, No. 3 3 Rotary arlers, No. 1 2 1st stage Cu-As purificatlon tank. 90 vol. tons, No. 3 Solution heaters, No. 3 Filter press, 30x30 bronze, No. 4 Zinc dust purification tanks, 45 vol tons. No. 3 5 4 Filter press, 36x36 bronze, No. 3 5 3 Cadmium recovery plant: Process tank. No. 5 2 Cake roaster, 20 ft diam x 4 hearth. No. 1 Filter press, 24x24, No. 4 1 Sponge wash box, 4x6 ft, No. 1 Evaporative cooling unit (vacuum), No. 1 Purified storage tank, vol. tons 400 400 400 Cell acid storage, vol. tons 400 400 400 Electrolytic cells, No. 180 372 372 Cell room ventilation, cu ft per min 35,000 125,000 125,000 Cell cooling water. gal per min 1,500 2,300 2,300 Deep well 16 in. x 95 ft, 1500 gal per min, No. 2 3 3 Melting and casting furnace, 130 ton. No. I Furnace fume scrubber unit, No. 1 Dross drums, No. 2 Dross roaster, 8 ft diam x 8 hearth, No. 1 Electrolysis power conversion, kw 6,250 23,750 23,750 Power transformers, 13,800/440, kva 1,000 1,500 2,000 Steam boilers fire tube, 15 psi, lb per hr 12,000 18,000 18,000 Steam boilers water tube. 125 psi, Ib per hr 30,000 Air compressor, 2 stage. 300 cu ft per min. 100 psi, No. 1 2 Air compressor, 1 stage, 300 cu ft per min, 20 psi, No. 1 Vacuum pumps, 18x7, 720 cu ft per day, No. - - . Vacuum pumps, 24x11, 1.633 cu ft per day, No. 3 3 Building area, sq ft 60,854 113,568 115,000 per ton of metal produced. In the summer of 1944, the "reverse" leaching process was placed in operation and since it has been described,' no further description will be given. Other facilities and changes which have contributed to the process improvements were the scrubbing of fume from the melting and

Jan 1, 1953

By R. L. Fleischer, W. F. Hosford

Tensile data for coarse grained aluminum Polycrystals suggest that the "grain size" effect is not due to dislocations piled up at grain boundaries but rather is primarily a relative size effect due to surface crystals being weaker and less confined. STUDIES directed at interpreting hardening of poly-crystalline metals normally identify their strain hardening properties with those in some particular type of single crystal.1"4 The recent recognition in face centered-cubic metals of a nearly linear stage with rapid hardening occuring at comparable rates for both polycrystals and single crystals, suggested that the same process or processes determine both cases and hence that there exists some justification for the use of single crystals to understand polycrystals. Further evidence for the above view may be found by an approach initiated by Chalmers:5 By using bicrystals of controlled orientation it is possible to begin to assemble a polycrystal and also to study grain boundary effects in detail. In this way it has been found that a single grain boundary affects easy glide but not the subsequent stage II hardening.6 This result suggests that a sensitive way to observe grain boundary effects in polycrystals would be to vary grain size and measure easy glide. As will be seen, easy glide is only possible for coarse-grained samples, and hence the results will serve to fill in the gap in measurements between single crystals and bicrystals on one hand and fine-grained polycrystals on the other. One problem inherent in comparing single crystals with polycrystals is the uncertainty as to what slip systems are acting in a polycrystal. To compare the two types of samples, rates of shear hardeninn---L. on the acting -planes are needed. and these may be computed only if it is known what particular systems are active. The acting systems were examined for a coarse-grained polycrystal and it will be shown that the systems supplying the preponderance of slip can be determined with little ambiguity. EXPERIMENTAL PROCEDURE Twelve samples of aluminum were prepared by chill casting into a heated graphite mold, followed by annealing at 635° ± 5°C for 24 hr with an 8-hr fur- nace cool, and finally either etching7 or electropol-ishing.' The samples, with a 7 to 10 cm length between grips and 4.4 by 6.6 mm in cross section, were deformed at a strain rate of about 3 10 -3 . per min in a tensile device which has been described elsewhere.5 The composition was reported by Alcoa as 99.992 pct Al, 0.004 pct Zn, 0.002 pct Cu, 0.001 pct Fe, and 0.001 pct Si; nine samples were deformed while immersed in liquid helium and three in air at room temperature. The stress-strain curve for one of the samples (P-1) deformed at 4.2 "K has been reported previ~usl~.~ This sample was selected for determination of active slip systems. Eighteen of the crystals were examined by optical microscopy to determine the angles of slip line traces and by X-ray back reflection to determine orientation. By this means the slip planes were determined and the resolved shear stress factors for possible slip systems could be computed. Finally each sample was sectioned so that after etching, the number of crystals could be counted for each of ten newly exposed surfaces. The average of these ten values will be termed n, the number of crystals per cross section. Values of 11, varied from 1.9 (nearly bamboo structure) to 12.7. Sketches of typical cross sections appear in Fig. 1. RESULTS AND DISCUSSION: SLIP SYSTEMS 1) Determination of Acting Slip Planes—The stress axis orientation and operative slip planes in eighteen crystals of sample P-1, as determined by slip line traces and crystal orientation, are summarized in Fig. 2. For one of the crystals two planes had a common trace. so that the traces alone did not distinguish which plane or planes were slipping. However it was found that the stress resolving factor for the primary system was 0.386, .while that for the most stressed system in the other plane (indicated bv the dotted arrow) is 0.138. It will be assumed tgerefore that only the primary plane acted. Since the orientations were determined after extending the samples 4 pct, the stress axes may be rotated from their original value by as much as 2 deg in some cases. It is interesting to note that in five crystals only one slip plane acted, in eight two acted, and in five three planes were observed—an average of two slip

Jan 1, 1962

By A. J. Lena, M. F. Hawkes

Changes in hardness, tensile properties, microstructure, electrical resistance, and X-ray diffraction effects indicate that lattice strains are necessary for the embrittlement of ferritic stainless steels when heated for relatively short times at 475°C (885°F). It is suggested that 475°C (885°F) embrittlement is due to the accelerated formation of an intermediate stage in the formation of s under the influence of these strains. FERRITIC stainless steels (low carbon alloys of iron with more than 15 pct Cr) are subject to two forms of embrittlement when heated in the temperature range of 375° to 750°C. The embrittlement which occurs after long time heating between 565" and 750°C is well understood; it is caused by the precipitation of the hard, brittle s phase. Sigma is an intermetallic compound of approximate equi-atomic composition with an extended range of formation in Fe-Cr alloys. The maximum temperature at which this form of embrittlement can occur is dependent upon chromium content; and is approximately 620°C for a 17 pct Cr steel and 730°C for a 27 pct Cr steel. The other form of embrittlement occurs after relatively short heating periods in the range of 375" to 565°C; in the higher chromium steels, hours may be sufficient as compared to months for s embrittlement. This phenomenon is not at all well understood and several controversial theories have been proposed. The rate and intensity of embrittlement increase with increasing chromium content but the maximum rate occurs at 475°C re-gardless of chromium content. As a result of this, the phenomenon has been termed 475°C (885°F) embrittlement. The effect of 475°C embrittlement on the properties of ferritic stainless steels has been thoroughly reviewed by Heger.1 The embrittlement causes a pronounced decrease in room temperature impact strength and ductility, a large increase in hardness and tensile strength, and a decrease in electrical resistivity and corrosion resistance. Microstructural changes accompanying embrittlement are minor and difficult to interpret with a general grain darkening, appearance of a lamellar precipitate, grain boundary widening, and precipitation along ferrite veins having been reported at various times. With the exception of reported line broadening, X-ray diffraction studies by conventional Debye analysis of solid samples have been of little value. BY making use of electron diffraction methods, Fisher, Dulis, and Car-roll' have recently shown the existence of a chromi-um-rich, body-centered cubic phase in 27 pct Cr steels which had been aged at 482°C (900°F) for as long as four years. Two types of theories have been advanced to account for the embrittlement. The first of these requires the precipitation of a phase not inherent in the Fe-Cr system with various investigators suggesting a carbide,3 nitride,3 phosphide,4 or oxide." Theories of this type have difficulty accounting for the influence of alloying elements on the embrittlement and for the facts that a minimum chromium content is necessary for embrittlement and the intensity of embrittlement increases with increasing chromium content. The second type of theory that has been proposed relates 475°C embrittlement to s phase formation which is inherent in the Fe-Cr system. An assumption of this kind can adequately explain the influence of alloying elements, for they exert an effect on 475°C embrittlement similar to that on s phase for-mation as can be seen in Table I. The minimum chromium content is essentially the same for both phenomena and it has been shown12,13 that s is a stable phase in the embrittling temperature range. In addition, it has been reported14,15 that pure alloys embrittle to the same extent as commercial type alloys. There are, however, several factors which have prevented complete acceptance of a s phase theory. Foremost of these is that the embrittlement can be removed by reheating for short time periods above 600°C, which in the higher chromium steels is within the stable s region. No s has ever been observed after one of these curing treatments, nor has any s been found as a result of embrittlement at 475°C. In addition, the simple precipitation of s cannot explain the time-temperature relationships for reactions between 350°and 750°C. This behavior is shown schematically in Fig. 1. Newell 16 and Ried-rich and Loib4 have shown that 475°C embrittlement follows a C-type curve as illustrated, while Short-

Jan 1, 1955

By N. J. Olson, G. W. Toop

A series of equations previously derived for calculating ternary thermodynamic properties using binary data has been applied to the problem of predicting ternary phase diagrams and quaternary excess free energy. The methods are considered to be rigorous for regular ternary and quaternary systerns and empirical for nonregular systems. The equations have been used to predict ternary phase boundaries in the Pb-Sn-Zn system at 926°K and the Ag-Pd-Cu system at 1000ºK. Calculated quaternary excess free-energy values are presented for the Pb-Sn-Cd-Bi system at 773°K. A method for predicting the location of ternary phase boundaries would be a useful supplement to experimental measurements in ternary systems. This has been recognized with the considerable work that has been done to find models to predict or extend thermodynamic properties and phase diagrams in binary and ternary systems1-18 for which direct experimental measurements are limited. With the access to highspeed digital computers and mechanical plotting devices, it is currently rather easy to compare mathematical models with experimental data. The regular-solution model is consistent with systems which exhibit negative heats of mixing, positive heats of mixing, and miscibility gaps, and therefore it is applicable to simple phase diagrams. The purpose of this paper is to illustrate the use of regular-solution equations to predict, empirically, phase equilibria in some types of nonregular ternary systems. Corresponding equations for regular quaternary systems are given and used to calculate empirical quaternary excess free-energy data. METHOD FOR PREDICTING THE LOCATION OF TERNARY PHASE BOUNDARIES USING BINARY DATA Meijerin1,6 has used the regular-solution model to calculate common tangent points to ternary free energy of mixing surfaces and hence to determine phase boundaries in ternary systems involving miscibility gaps. He used the following equation to calculate ternary excess free energy of mixing values: stants characteristic of the binary solutions, and Ni is the mole fraction of component i. An alternate expression which gives for regular solutions as a function of binary values of along composition paths with constant N1/N2, N2lN3, and N1/N3 may also be derived:15 ternary r xs 1 ?c-*n.Ti*.U*. This expression for is more useful for the empirical calculation of ternary excess free-energy values for nonregular systems because actual binary AFXS data may be used in the expression rather than attempting to find suitable constants for Eq. [I]. The results of this feature of Eq. [2] are illustrated in Table I where calculated excess free-energy values for the Ni-Mn-Fe system at 1232°K are compared with experimental data of Smith, Paxton, and McCabe.19 Although regular-solution equations have been shown to give calculated thermodynamic quantities which agree quite well with experiment for single-phase nonregular ternary systems,14,15 care should be exercised in the use of the equations to predict thermo-dynamic properties of multiphase ternary systems in which strong compound formation is suspected. This precaution is consistent with the simple regular-solution model which for negative values of ai_j will indicate a tendency toward compound formation but even for very large negative values of ai-jwill not give multiphase binary or ternary systems involving a distinct stable compound. Hence, calculated ternary free-energy data using Eq. [2] might be expected to vary between being rigorous and poor, in the following order, for ternary systems which are: a) regular solutions, b) nonregular single-phase liquids in which random mixing is nearly realized, c) nonregular single-phase solids, d) nonregular multiphase systems exhibiting miscibility gaps, e) nonregular multiphase systems with binary compounds but no ternary compounds, f) nonregular multiphase systems with highly stable binary and ternary compounds. The calculated data will be expected to be least accurate for the last two cases. The general method adopted in this paper involves two-dimensional plots of ternary activity curves. The principle used is that tie lines indicating two-phase equilibria join conjugate phases a and B for example, for which a1(a) = a1(B), a2(a) = a2(B), and a3(a) = a3(B). Tie lines may be determined by plotting the ternary activities of two components along an isoactivity line for the third component and the unique points where the above equalities hold may be found graphically.

Jan 1, 1970

P. L. Pratt (University of Birmingham, Birmingham, England)—The author has measured the hardening effect of isolated edge and screw dislocation boundaries in a remarkably elegant manner, and he proposes that the detailed structure of the boundary determines its effectiveness as a barrier to slip although "there are still unknown details of the distribution which may be extremely important." On p. 678 the author states "that for a good crystal (one having a yield of about 30 psi) a change in annealing temperature from 300" to 400 °C resulted in a negligible shift in the level of the stress-strain curve." Is this in agreement with the results of the earlier paper by Li, Washburn, and Parker8 in which such a change of temperature produced notable hardening in crystals with a yield of about 30 psi? Or were these earlier crystals, of the same nominal purity, less perfect than those used in this work? The experiment illustrated in Fig. 15, p. 681, shows that the boundary has sharpened after annealing at 400 °C, the temperature at which the hardening effect first becomes apparent. Did the author verify that no sharpening of the boundary occurred at, say, 300 °C where no hardening was found? This would seem to be an important check on the proposed mechanism, especially since zinc of this purity should polygonize at temperatures well below 400 °C within the annealing times used here. There seem to be at least two alternative explanations which could fit these facts: 1—The boundary may sharpen at 300 °C or even lower temperatures but be unable to contribute greatly to the hardening until jogged by thermal vacancies and other impurity atoms after 400 °C anneal. This mechan- ism, discussed in the earlier paper,' is believed to account for the thermal hardening observed by Blank'" in NaCl, as stated elsewhere: and by results from recent experiments which tend to confirm this view. 2—The dislocations in the diffuse boundary may be locked by impurity atoms at 300 °C and thus only sharpen the boundary at 400 °C with the aid of thermal vacancies. In this case it would be difficult to distinguish between the two sources of hardening (sharpening or jogging), using metal of this purity. Jack Washburn (author's reply)—The author wishes to thank Dr. Pratt for his discussion. In connection with his first comment, it should be pointed out that the temperature of testing in the present series of experiments was —196°C, whereas, in the earlier experiments he referred to,' it was 20°C. Therefore, the yield values should not be directly compared. It is likely that the crystals used in the present work were initially more perfect macroscopically than the large shear specimens used previously. The small size of the test section in the kink specimens made it possible to select regions of the large crystals, from which they were acid-cut, that were free of observable small angle boundaries. The sharpening of a microscopically diffuse boundary such as that in Fig. 15 begins in zinc of this purity well below 300°C. However, it is probably not justified to conclude that because a boundary looks sharp under the microscope it has attained the ideal structure. The author agrees with Dr. Pratt that jogging of dislocation lines is one of the structural features of a small angle boundary that should be important in determining its strengthening effect. More detailed information concerning the changes in boundary structure as a

Jan 1, 1956

By C. E. Williams, G. H. Bruce

The trend toward deeper drilling, together with the attcndant increase in power requirements for circulation of the drilling fluid, has emphasized the need for a critical examination of the factors affecting the removal of bit cuttings from the hole by the drilling fluid. The ability of drilling fluids to lift cuttings is called their carrying capacity. A series of laboratory and field experiments has been conducted to determine the minimrim annular velocity necessary to remove cuttings, and to investigate the effects of properties of drilling fluids on their carrying capacities. Consideration of the results of these experiments led to the following conclusions: 1. Turbulent flow in the well annulus is most desirable from the standpoint of cutting removal. 2. Low viscosity and low gel are advantageous in removing cuttings. 3. Increase in mud weight is effective in increasing carrying capacity. 4. The carrying capacity is higher when the pipe is rotated than when it is not. 5. If turbulent flow can be maintained, an annular velocity slightly higher than the slip velocity of the largest cuttings to be transported should keep the bore hole clean. This implies velocities of 100 to 125 ft per minute rather than the presently used 175 to 225 ft per minute. INTRODUCTION Power Savings by Reduction of Annular Velocities A large portion of the power expended in drilling operations is consumed in circulating the drilling fluid. An important factor in establishing the rate of mud circulation is the minimum velocity in the annulus necessary to remove bit cuttings. Empirically, it has been found that average annular mud velocities of about 200 ft per minute will remove cuttings. It was not definitely known, however, whether annular velocities of about 200 ft per minute were just above the minimum necessary to remove cuttings, or whether such velocities could be materially reduced without sacrifice of the ability of the mud to remove cuttings. It is apparent that if annular velocities could be reduced without impairment of cutting removal, a considerable saving in power requirements would result. Need for Research on Carrying Capacity The ability of a drilling fluid to transport cuttings is called its carrying capacity. Although it has been recognized that the carrying capacity of a mud is affected by mud properties such as viscosity and density1,2,3,4 there have been various views in the industry as to the effects of these mud properties on carrying capacity. The economic importance of the problem of carrying capacity and the scarcity of information on the subject indicated that research on the problem was needed. THE FACTORS AFFECTING CARRYING CAPACITY Qualitative Determination of Factors The mechanism of cutting transport is closely related to that involved in the separation of material by settling processes. A considerable amount of research has been done on settling problems, and discussions of sedimentation theory can be found in standard texts.5,6 Consideration of the information available from these sources, together with consideration of the mechanism of cutting transport, leads to the conclusion that the factors affecting carrying capacity are the dimensions of the system, the physical properties of the cuttings. and the physical properties of the drilling fluid. system Dimensions and Their Effect The dimensions of the fluid circulating system of importance to carrying capacity are the bore hole size, drill pipe size, pump capacity, and pump speed. These dimensions determine the annular velocity of the drilling fluid. Physical Properties and Their Effects The physical properties involved in the interaction between mud and cuttings are the density and shape of the cuttings and the density, viscosity and gel strength of the drilling fluid. The effect of the density factor on carrying capacity is fairly obvious; high density difference between cuttings and fluid results in a low buoyant force and therefore decreases carrying capacity. The effect of cutting shape is less obvious. Although

Jan 1, 1951

By R. J. Tailleur, M. Rosenberg

Oil well drilling muds prepared with lubricating aids such as oil, graphite or mica will not produce a sufficiently strong protective lubricant film for the bit bearing surfaces under high load conditions encountered in drilling operations. Consequently, the damage to the bearing surfaces resulting from the absence of such a film frequently results in a shortening of the drill bit life. Several different types of extreme pressure lubricating additives for use in muds have been developed. These arlrlitives are capable of providing a protective lubricating film for bearing surfaces subjected to high load pressures. A modified version of the Timken lubricant tester, an instrument commonly used for testing the extreme pressure lubricating properties of oils, is used for measuring and controlling the lubricating qualities of the mud. Comparison tests made in the laboratory with 55/8-in. tricone bits drilling on steel plates immersed in extreme pressure lubricating muds and conventional muds showed that the bearing life of those bits tested in the lrrbricating mud was three to four times greater than obtained with conventional muds. Drilling tests made with extreme pressure lubricating muds in West Texas increased over-all bit life three to fourfold, reduced drilling torque and wear on drill pipe and drill collars. INTRODUCTION Roller bit cones contain roller, ball and journal-type bearings that are subjected to very heavy loading during drilling operations in which the only lubricant available for these bearings is either water or drilling mud. These bearings must be able to operate satisfac- torily in order for the cutting structures on the surface of the cones to be fully utilized. Frequently, the bearings wear to the point that the cones freeze and the bit must be pulled before the cutting structures on the cones become worn. In some instances the bearings continue to turn with severe wear but the cones, which are held on the bit by the bearings, drop off; and this causes an expensive fishing job. In other instances the bearings wear to the point that the cones pinch in causing an undergauge hole and delays to ream the hole. The advantage of increasing bit bearing life is self-evident. Because of inability to control the internal wear of bits while drilling, it was felt that improvement of the lubricating qualities of the circulating mud would be helpful. However, the often-discussed lubricating property attributed to muds has been a rather vague and intangible one. The literature has implied that the oil-emulsion muds are more effective lubricants than non-emulsion muds. Actually, there has been no direct method of measuring the lubricating property of muds which would indicate that one mud was better than another; there has been no way to determine whether the lubricating property of any mud was adequate. Furthermore, in comparing lubricating properties of muds, two qualities should be considered—the coefficient of friction and the extreme pressure lubricating qualities. For example, graphite can reduce the coefficient of friction, but does not provide a strong enough lubricating film to protect bit bearings subjected to high loads. The primary object of the investigation reported herein was to find a method for prolonging the bearing life of drill bits; and considering the heavy loads used in drilling, it was felt necessary to find both a method of measuring lubricating properties under high load conditions and methods of establishing effective lubrication under these conditions. Extreme pressure lubrication deals with that realm of lubrication where metal surfaces are in rubbing contact with each other under very high pressures. These pressures may be quite high, as for example,